Silicon carbide possesses properties that should make it a superior semiconductor for applications that involve high temperature, high power, high radiation and/or high frequency (See Amorphous and Crystalline Silicon Carbide and Amorphous and Crystalline Silicon Carbide II). In addition, a variety of optical devices (e.g. light-emitting diodes, LEDs) can be fabricated from SiC and operated at high temperature. The properties that allow this superior performance are its large bandgap, excellent physical stability, high thermal conductivity, high electric breakdown field, and high saturated electron drift velocity. Semiconductor devices fabricated from SiC are capable of operating at temperatures above 600.degree. C., which is well above the capability of current commercial semiconductors. Also, based on theoretical considerations, SiC microwave devices should be far superior to currently available devices. The potential market for SiC devices is very large and could exceed billions of dollars annually.
The chief obstacle to the commercialization of SiC has been the lack of control over its crystal growth. Several properties of SiC contribute to this lack of control. First, it does not melt at reasonable pressures; it sublimes at temperatures above 1800.degree. C. Second, it grows in many different crystal structures, called polytypes. The following defines some of the nomenclature used in this document. The SiC polytypes are formed by the stacking of double layers of Si and C atoms. Each double layer may be situated in one of three positions, designated as A, B, and C. The sequence of stacking determines the particular polytype and the stacking direction is called the crystal c-axis. There is one cubic polytype, with the zincblende structure and known as 3C or .beta. SiC. It has a three-layer repeat sequence ABC . . . , and so forth. All of the other polytypes are known as .alpha.-SiC and have either a hexagonal or rhombohedral structure. The hexagonal 6H-SiC polytype has the six layer repeat sequence ABCACB . . . , and so forth. For the .alpha.-SiC polytypes, the (0001) plane is known as the basal plane, and this plane is perpendicular to the c-axis. The (111) plane of 3C-SiC is equivalent to the basal plane of the .alpha.-SiC polytypes. In the discussions of this document, "basal plane" shall refer to either the (0001) plane for .alpha.-SiC, or the (111) plane of 3C-SiC. Also, "3C" shall be used for 3C-SiC, and "6H" shall be used for 6H-SiC. It is well known that a SiC surface, which is approximately parallel to the basal plane, is terminated with either Si atoms (the surface called the Si face) or terminated with C atoms (the surface called the C face). The term "vicinal (0001) wafer" shall refer to SiC wafers whose polished surface is misoriented less than 6.degree. from the basal plane. The term "homoepitaxial" shall refer to epitaxial growth whereby the film and the substrate (wafer) are of the same polytype, and the term "heteroepitaxial" shall refer to growth whereby the film is of a different polytype than the substrate.
Since melt-growth techniques cannot be applied to SiC, vapor growth processes have been developed. An early SiC vapor-growth process, the Lely process (J. A. Lely, Ber. Dt. Keram. Ges. 32, 229 (1955)), was based on the sublimation of polycrystalline SiC within a growth cavity and did produce rather pure crystals of various polytypes. Unfortunately, the crystals were too small, irregular in shape, and thus not suitable for commercial development. More recently, a high temperature sublimation process has been developed that does produce large single-crystal boules of 6H. Furthermore, polished wafers, more than 25 mm in diameter, can be produced from these boules. Commercial 6H-SiC devices are now being produced with these wafers.
In order to fabricate semiconductor devices, it is necessary to be able to grow high-quality, low-defect-density, single-crystal films on single-crystal wafers. A variety of processes have been developed to produce homoepitaxial SiC films at relatively high temperatures (above 1600.degree. C.). It is well known that single-crystal homoepitaxial 6H films can be grown in the lower temperature range 1400.degree. C. to 1550.degree. C. by chemical vapor deposition (CVD) on vicinal (0001) 6H wafers if the tilt angle is greater than about 1.5.degree.. Typically 3.degree. to 4.degree. is used. Large tilt angles can cause problems. Since .alpha.-SiC has anisotropic properties, CVD films, grown on vicinal (0001) substrates with non-zero tilt angles, will have anisotropic properties parallel to the growth surface. The larger the tilt angle, the larger will be the anisotropy. This could result in unwanted behavior in devices fabricated with the CVD films. From this viewpoint, small tilt angles are desirable.
If the tilt angle of a vicinal (0001) 6H-SiC substrates is less than about 1.degree., a 3C heteroepitaxial film will be produced on the 6H wafer if prior art processes are used. However, the 3C film grown in this way generally has a high density of defects, including a defect known as double positioning boundaries (DPBs). This DPB defect can arise because of the change in stacking sequence of the 6H wafer (i.e. ABCACB . . .) to that of the 3C (ABC. . . or ACB. . .) film at the interface between the two polytypes. The difference between the two 3C sequences is a 60.degree. rotation about the &lt;111&gt; axis. If both of these sequences nucleate on the 6H substrate, DPBs will form at the boundary between domains differing by the 60.degree..
A theoretical crystal growth model, proposed by Matsunami and which has been used to explain the formation of the 3C and 6H polytypes on vicinal (0001) 6H substrates, is based on the density of atomic-scale steps on the growth surface. According to this model, 6H grows on 6H when the tilt angle is greater than about 1.5.degree. because terraces between steps are small and arriving molecules containing Si and C are able to migrate to steps where growth occurs. This growth is a lateral growth of the steps and reproduces the 6H substrate. For small tilt angles, say less than 1.degree., the terraces are larger and all molecules are not able to migrate to steps; instead, nucleation of 3C takes place on the terraces. Hence, for small tilt angles, 3C grows on 6H substrates. Thus, for the case of growth on low-tilt-angle (less than 1.degree.) polished 6H wafers, prior art teaches that a 3C film is inevitable.
In the CVD growth of epitaxial SiC films on vicinal (0001) .alpha.-SiC wafers, a variety of pregrowth process have been used to prepare the polished surface for growth. The intent of these processes is to remove contamination and near-surface defects (from cutting and polishing the wafer) that contribute to poor quality films. Processes used to eliminate near-surface defects include molten-salt etching, oxidation followed by removal of the oxide with hydrofluoric acid, reactive ion etching, etc. prior to loading wafers into the CVD growth system. In situ processes (within the CVD growth system) include high temperature etching in H.sub.2 or HCl/H.sub.2 mixtures. All prior art use of these processes in CVD at temperatures less than 1600.degree. C. has produced the heteroepitaxial growth of 3C on vicinal (0001) 6H if the tilt angle was less than 1.degree.. In one case, Powell et al. used an HCl etch consisting of 2 min at 1200.degree. C. In another case, Matsunami et al. used an HCl etch consisting of 10 min at 1500.degree. C. In both of these cases, 3C films were produced on vicinal (0001) 6H with tilt angles of less than 1.degree.. It must be emphasized at this point that prior-art teaches that pregrowth surface treatments of substrates used in CVD processes, can be effective in reducing defects in the resulting film. Prior art has not taught that surface treatments can be a significant factor in controlling the polytype of CVD-grown SiC films.